JP7263696B2 - Alloys for rare earth magnets - Google Patents

Description

The present invention relates to an alloy for rare earth magnets.

In recent years, there has been a demand for the development of magnets with a reduced content of rare earth elements. As used herein, the rare earth element refers to at least one element selected from the group consisting of scandium (Sc), yttrium (Y), and lanthanides. Here, the lanthanoid is a general term for 15 elements from lanthanum (La) to lutetium (Lu). RT ₁₂ (R is at least one rare earth element, T is Fe, Co or Ni) having a body-centered tetragonal ThMn ₁₂ type crystal structure is a ferromagnetic alloy containing a relatively small composition ratio of rare earth elements. Are known. Although RT ₁₂ has high magnetization, it suffers from thermal instability of the crystal structure.

Patent Document 1 discloses a rare earth permanent magnet in which a portion of Fe, which is a T element, is partially replaced with Ti, which is a structure stabilizing element, to improve thermal stability in exchange for high magnetization. there is

In Patent Document 2, by partially substituting the R element of the RFe ₁₂ -based compound with the substituting element M1 such as Zr and Hf, the amount of the substituting element M2 such as Ti substituting the transition metal element is reduced and saturated. Rare earth permanent magnets with stabilized _ThMn12 structures are disclosed while retaining their magnetization.

Further, Patent Document 3 discloses an R'--Fe--Co system ferromagnetic alloy in which Y or Gd is selected as part of the R element of RFe ₁₂ , and this R'--Fe--Co system ferromagnetic It is described that the alloy exhibits high magnetic properties due to having a ThMn ₁₂ -type crystal structure produced by an ultraquenching method.

Further, Patent Document 4 describes that by adding Cu, a non-magnetic and low melting point 1-4 composition (SmCu ₄ phase) phase is generated, and sintering and high coercive force are possible. .

Further, in Patent Document 5, at least one of Sm ₅ Fe ₁₇ -based phase, SmCo ₅ -based phase, Sm ₂ O ₃ -based phase, and Sm ₇ Cu _3- based phase is added as a secondary phase to a ThMn _12- type main phase. It is described that the coercive force can be increased by including the element.

Further, Patent Document 6 describes that the addition of Cu generates a liquid phase and enables the formation of a dense bulk body.

JP-A-64-76703 JP-A-4-322406 JP 2015-156436 A Japanese Patent Application Laid-Open No. 2001-189206 Japanese Unexamined Patent Application Publication No. 2017-112300 WO2016/162990

The single-crystal-like main phase particles used in anisotropic magnets, which are often used as high-performance magnets, are obtained by efficiently pulverizing the raw material alloy (material to be pulverized) into single-crystal units during fine pulverization. . Furthermore, considering the general treatment temperature during the sintering process, the main phase compound is also required to exist stably at least at 900° C. or higher, preferably at 1000° C. or higher.

The rare earth permanent magnet described in Patent Document 1 has improved thermal stability due to element substitution of Fe with Ti, but since the amount of Fe substituted with Ti is large, the magnetization is correspondingly reduced, and sufficient magnetic properties are obtained. I can't get it.

On the other hand, in the rare earth permanent magnet described in Patent Document 2, the _ThMn12 structure is stabilized by substituting a transition metal element with Ti or the like, but the effect is not necessarily sufficient.

The R'--Fe--Co ferromagnetic alloy described in Patent Document 3 does not replace the Fe element with the structure stabilizing element M (such as Ti), so it exhibits high magnetization, large magnetic anisotropy, and a high Curie temperature. However, due to the non-equilibrium phase, the main phase compound may decompose in a high temperature densification process such as sintering.

The rare earth magnet described in Patent Document 4 may not have high magnetic physical properties due to the large amount of Ti added.

In the rare earth magnet described in Patent Document 5, when the rare earth-rich secondary phase Sm ₇ Cu ₃ is used, there is concern that the equilibrium state will shift to a composition richer in rare earth than the main phase during heat treatment, and the main phase ratio will decrease. be done.

In the rare earth magnet described in Patent Document 6, since the Fe element is not replaced with the structure stabilizing element M, high magnetization, large magnetic anisotropy, and a high Curie temperature can be obtained, and the density as a bulk body is high. Since it is a non-equilibrium phase, the main phase compound may decompose in a high temperature process such as sintering at 1000° C. or higher. Furthermore, it is difficult to obtain an anisotropic sintered magnetic powder.

Accordingly, an object of the present disclosure is to provide a rare earth magnet alloy suitable for obtaining anisotropic sintered magnetic powder.

The rare earth magnet alloy of the present disclosure, in a non-limiting exemplary embodiment, is a rare earth magnet alloy having a primary phase and one or more secondary phases, wherein the composition of the entire alloy is the following composition formula (1 ).
R(Fe,Co) _{w- zTizCuα (} 1)
Here, R is at least one rare earth element, and w, z, and α are 8≤w≤13, 0.42≤z<0.70, and 0.40≤α≤0.70, respectively. Be satisfied.

In one embodiment, the overall composition is represented by the following compositional formula (2).
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α (2)
Here, R is composed of R1 and R2, R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1, and R2 is selected from the group consisting of Sm, La, Ce, Nd and Pr and must contain Sm, and Sm is 50 mol% or more of the entire R2. x and y respectively satisfy 0.5≦x≦1.0 and 0≦y≦0.4.

In one embodiment, the main phase has a ThMn ₁₂ -type crystal structure, and the composition of the main phase is represented by the following compositional formula (3).
R1 _1-x' R2 _x' (Fe _1-y' Co _y' ) _12-z'-α' Ti _z' Cu _α' (3)
where x′, y′, z′ and α′ are respectively 0.5≦x′≦1.0, 0≦y′≦0.4, 0.48≦z′<0.91, and 0.24≦α′≦0.37.

In some embodiments, z' satisfies 0.48≤z'<0.74.

In one embodiment, the main phase is a phase having a ThMn _12- type crystal structure, and the secondary phase is mainly a crystalline phase in which 50 mol % or more of the entire secondary phase is composed of Cu.

In one embodiment, the crystal structure of said subphase is KHg ₂ type.

In one embodiment, the KHg type ₂ crystal structure phase is 50% or more by volume of said minor phase.

In one embodiment, the subphase comprises R atoms and the R atoms present in the subphase have a molar ratio of [R2]/([R1]+[R2]) higher than the composition of the entire alloy.

In one embodiment, the weight ratio of the secondary phase in the entire rare earth magnet alloy is 3 wt % or more and 10 wt % or less.

In one embodiment, the subphase absorbs and releases hydrogen.

In some embodiments, hydrogen absorption and desorption occurs at temperatures of at least 700° C. or less.

In one embodiment, cracks are formed between the main phase and the subphase or in the subphase.

In one embodiment, the maximum value of the X-ray intensity in a powder X-ray diffraction pattern with a lattice spacing value of 3.03 Å or more and 3.10 Å or less is 2.50 Å or more and 2.57 Å. The X-ray intensity is greater than the maximum value in the following range, and the α-(Fe, Co, Ti) phase is 10% by weight or less.

In one embodiment, the α-(Fe,Co,Ti) phase is 5 wt% or less.

According to the embodiments of the present invention, anisotropic sintered magnetic powder with improved magnetic properties and thermal stability can be obtained with high efficiency.

For the rare earth _magnet alloy having _the composition of _Y0.4Sm0.6 ( _Fe0.80Co0.20 ) _{8.79Ti0.42Cu0.67} of Example 6 in the embodiment of the present disclosure, 1050 _{° C} for 20 minutes 1 is a diagram showing the results of observation of a cross-sectional structure with a polarizing microscope before heat treatment of . For the rare earth magnet alloy having the composition of Y _0.4 Sm _0.6 (Fe _0.80 Co _0.20 ) _{8.79 Ti0.42} Cu _0.67 of Example 6 in the embodiment of the present disclosure, It is a diagram showing the observation results of the cross-sectional structure with a polarizing microscope after the heat treatment of . FIG. 1C is a backscattered electron (BSE) image obtained with a scanning electron microscope (SEM) of the sample of FIG. 1B; FIG. 1B shows the results of compositional analysis of each phase in the sample of FIG. 1B; FIG. 5 is a diagram showing temperature dependence of volume magnetization of main phases of Examples 3 to 6 in the embodiment of the present disclosure; 4 is a graph showing the temperature dependence of magnetic anisotropy fields of samples of Examples 3 to 6 in the embodiment of the present disclosure; FIG. 10 shows powder X-ray diffraction patterns before and after heat treatment at 1050° C. for 20 minutes of the sample of Example 10 in the embodiment of the present disclosure; FIG. 13 is a diagram showing the temperature dependence of the hydrogen discharge amount of Example 13 in the embodiment of the present disclosure; FIG. 11 shows powder X-ray diffraction patterns of samples of Examples 15-18 in embodiments of the present disclosure. FIG. 11 shows a BSE image obtained by SEM of the sample of Example 15 in an embodiment of the present disclosure; FIG. 12 shows a BSE image obtained by SEM of the sample of Example 16 in an embodiment of the present disclosure; FIG. 11 shows a BSE image obtained by SEM of the sample of Example 17 in an embodiment of the present disclosure; FIG. 11 shows a BSE image obtained by SEM of the sample of Example 18 in an embodiment of the present disclosure;

[Composition of alloy for rare earth magnet]
The rare earth magnet alloy of the present disclosure, in an exemplary, non-limiting embodiment, has a primary phase and a secondary phase, and the overall composition is represented by compositional formula (1) below.
R(Fe,Co) _{w- zTizCuα (} 1)
Here, R is at least one rare earth element. Also, w, z, and α satisfy 8≦w≦13, 0.42≦z<0.70, and 0.40≦α≦0.70, respectively.

In one embodiment, the overall composition is represented by the following compositional formula (2).
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α (2)
Here, R is composed of R1 and R2. R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always containing Sm, Sm is 50 mol% or more of the entire R2. R1 is preferably Y only (excluding unavoidable impurities), and R2 is preferably Sm only (excluding unavoidable impurities). x and y respectively satisfy 0.5≦x≦1.0 and 0≦y≦0.4.

By including Sm in the rare earth element R, the magnetic anisotropy of the main phase, which is important for increasing the coercive force, can be improved.

As a result of diligent research by the present inventors, as shown in the above compositional formulas (1) and (2), by adding Cu to the raw material alloy, the melt of the raw material alloy is rapidly cooled and solidified into the alloy. It was found that a rare earth-rich phase (subphase) coexisting with the main phase (hard magnetic phase with high magnetization and magnetic anisotropy) was formed. It was found that the formation of the secondary phase richer in rare earth elements than the main phase facilitates crystal growth of the main phase by heat treatment performed on the rapidly solidified alloy. This heat treatment also makes it possible to easily reduce heterogeneous phases during melting and solidification of the material alloy. Furthermore, the absorption and release of hydrogen by the rare earth-rich subphase causes cracks to occur between the main phase and the subphase or in the subphase, enabling efficient pulverization into single crystal units. These are extremely beneficial in obtaining anisotropic sintered magnetic powder, and can enable mass production of highly oriented anisotropic sintered magnetic powder.

The amounts of R1, R2 and Ti affect the magnetic properties and high temperature stability of the main phase. From the viewpoint of magnetic anisotropy, R2 is desirably half or more of R1 (half or more of the total of R), and the preferred range of x is 0.5≤x≤1.0. From the viewpoint of saturation magnetization, Ti should preferably be as small as possible, but from the viewpoint of high-temperature stability, it should be large. A range of 0.42≦z<0.70 is suitable. In particular, when z is 0.70 or more, the saturation magnetization, magnetic anisotropic magnetic field and Curie temperature all decrease. Note that 50 mol % or less of Ti may be replaced with tungsten (W), vanadium (V), or the like.

From the viewpoint of increasing the magnetic moment and improving the magnetization and magnetic anisotropy at practical temperatures associated with the increase in the Curie temperature, it is preferable to replace part of Fe with Co. However, if the amount of substitution with Co is too large, the magnetization and magnetic anisotropy will rather be lowered. Specifically, the Co substitution amount y is desirably 0≤y≤0.4, and more desirably 0.1≤y≤0.3.

The amount of Cu is set so that the amount of the generated subphase is an appropriate value. If the amount of the subphase is small, not only is it impossible to eliminate the heterogeneous phase during melting and solidification of the raw material alloy, but it is also difficult to grow crystals to a size sufficient to obtain anisotropic sintered magnetic powder. Also, when the amount of the sub-phase is large, the ratio of the main phase is lowered, so the magnetization of the magnet body is lowered. According to the inventor's experiments, the appropriate amount of Cu is in the range of 0.40≦α≦0.70.

50 mol % or more of the entire secondary phase is a crystal phase having a Cu composition (that is, a Cu group richer in R than the main phase). In some embodiments, the minor phase comprises predominantly a KHg type ₂ crystal structure phase (hereinafter 1-2 phase). The subphase may also include a phase having a composition in which the ratio of "R" to "Cu, Fe, and/or Co" is 1:4 (hereinafter referred to as 1-4 composition). Regarding the R element constituting the subphase, both phases have a higher molar ratio of [R2]/([R1]+[R2]) than the composition of the entire alloy. In addition, Fe and Co may slightly dissolve in the subphase. Ti hardly dissolves in both phases.

An appropriate amount of w varies depending on the amount of Cu added to the raw material alloy, but 8≤w≤13. If w is too large, a soft magnetic α-(Fe, Co, Ti) phase is produced. If w is too small, 2-17 phases and 3-29 phases are formed. None of these phases are preferable for obtaining magnets with high magnetic properties.

The main phase of the alloy for rare earth magnets thus obtained has a ThMn type ₁₂ crystal structure in the embodiment. The TnMn ₁₂ -type compound phase in the alloy in the present disclosure can typically exist stably even at 1000° C. or higher. As such, the alloy embodiments of the present disclosure can be suitably used to employ high performance magnet fabrication processes such as sintering.

In general, the "ThMn ₁₂ -type crystal structure" is a tetragonal crystal, but in the present invention, the crystal lattice of the tetragonal crystal is slightly distorted to have orthorhombic symmetry, and the atoms in the crystal Even if the periodicity of is slightly disturbed, it is regarded as "ThMn type ₁₂ crystal structure".

Since the main phase to be generated contains Cu, the magnetic property values are different when compared with the main phase not containing Cu, even with the same amount of Ti substitution. First, the saturation magnetization decreases at least by the solid solution amount of Cu and Ti. The magnetic anisotropic magnetic field exhibits complex behavior due to the co-substitution of Cu and Ti. Specifically, contrary to the general characteristics of ThMn ₁₂ -type crystal structure compounds in which Cu is not substituted, the magnetic anisotropy field tends to decrease with the addition of Ti. Therefore, it is preferable that both Cu and Ti be as small as possible from the viewpoint of magnetic property values. However, from the viewpoint of high-temperature stability, a large amount of Ti is desirable. Specifically, when the main phase is represented by the composition formula R1 _1-x' R2 _x' (Fe _1-y' Co _y' ) _12-z'-α' Ti _z' Cu _α' , 0.5 ≤ x' ≤ 1.0, 0 ≤ y' ≤ 0.4, 0.48 ≤ z'< 0.91, 0.24 ≤ α' ≤ 0.37 are appropriate, and the substitution amount of Cu and Ti is More preferably, 0.48≦z′<0.74 and 0.24≦α′≦0.37.

[Method for producing alloy for rare earth magnet]
<Step A> Melting and solidifying step As a method for producing the R-Fe-Co-Ti-Cu-based rare earth magnet alloy, known methods such as mold casting, centrifugal casting, strip casting, and liquid ultra-quenching can be used. method can be adopted. These methods involve making a molten alloy and then cooling and solidifying the molten alloy. During solidification of the molten alloy, it is desirable to minimize the formation of phases (heterogeneous phases) such as α-(Fe, Co, Ti) phases, which are particularly unfavorable as raw material alloys for magnets. By adopting a method of supplying a molten metal onto a rotating roll and solidifying it to produce an alloy in the form of ribbons or flakes, such as a strip casting method or a liquid ultra-quenching method, which has a relatively high cooling rate, such a heterogeneous phase generation can be suppressed. When the cooling rate during solidification is low, the grain size of the precipitated heterogeneous phase becomes large. When the grain size of the heterogeneous phase contained in the alloy becomes large, it is difficult for the heterophase to disappear in the subsequent heat treatment step B.

When a molten alloy is quenched and solidified at a high cooling rate as in the liquid ultraquenching method, "nanocrystals" with nanometer-order sizes are produced in the solidified alloy. Anisotropic magnetic powder cannot be obtained by pulverizing the solidified alloy in the state of "nanocrystals". However, even nanocrystals can be easily grown into crystal grains of 10 μm or more suitable for obtaining anisotropic magnetic powder by going through the subsequent heat treatment step B.

<Step B> Heat Treatment Step By applying heat treatment to the alloy of the present invention, the following can be achieved.
(1) It reduces the amount of foreign phases produced during the solidification process.
(2) Coarse crystal grains. This is effective for easily obtaining a powder composed of single-crystal-like particles, which is useful as a raw material for an anisotropic sintered magnet, by a pulverization method.

The melting point of the 1-2 phase is around 860.degree. C., and the melting point of the 1-4 composition phase is around 880.degree. Therefore, the heat treatment temperature is preferably 900° C. or higher and 1250° C. or lower, more preferably 1000° C. or higher and 1100° C. or lower.

The heat treatment time depends on the heat treatment temperature, but is preferably from 5 minutes to 50 hours. If the time is too short, there will be insufficient reaction to eliminate the heterogeneous phase or insufficient grain growth. If the time is too long, vaporization and oxidation of the rare earth elements will occur, and the operational efficiency will be poor. At this heat treatment temperature, at least part of the subphase becomes a liquid phase, and part of the main phase dissolves and reprecipitates. For this reason, the crystal grains grow dramatically in the main phase as compared with the case where the liquid phase is not generated. In addition, heterogeneous phases generated during rapid solidification of the molten alloy can be easily eliminated if the grain size is not large.

The weight ratio of the subphase in the alloy before the step B is desirably 3 wt % or more and 10 wt % or less of the entire alloy. If the amount of the subphase is small, not only will the heterogeneous phase in the alloy not disappear even by the heat treatment in step B, but it will also be difficult to grow the crystals to a size sufficient to obtain an anisotropic sintered magnetic powder. Also, when the amount of the sub-phase is large, the ratio of the main phase is lowered, so the magnetization of the magnet body is lowered. When the actual amount of the subphase was identified by the X-ray Rietveld analysis of the sample with the weight ratio of the subphase of 3%, it was found that the amount of the subphase introduced was equivalent to the amount of the subphase in the alloy.

<Step C> Step of introducing cracks by absorption and release of hydrogen Since the subphase contained in the present alloy has a rare earth-rich composition containing a relatively large amount of rare earth elements, it can absorb and release hydrogen. Especially by heat treatment in hydrogen, the subphase can take up hydrogen significantly. With this alloy, for example, hydrogen absorption occurs at temperatures between 250°C and 400°C, and hydrogen release occurs between 300°C and 660°C. Therefore, after the alloy is heated to, for example, 350° C. in hydrogen to absorb hydrogen, the atmosphere can be switched to a vacuum atmosphere or an inert gas atmosphere to release hydrogen. In that case, the temperature for switching to the vacuum atmosphere or the inert gas atmosphere is preferably less than 600°C. When this alloy is exposed to a hydrogen atmosphere at a temperature of 600° C. or higher, there is a possibility that a different phase such as an α-(Fe, Co, Ti) phase is generated due to decomposition of the main phase due to a hydrogenation-disproportionation reaction. It is desirable that the ratio of α-(Fe, Co, Ti) phase in the alloy after hydrogen treatment is 10% by weight or less. More preferably, the proportion of α-(Fe, Co, Ti) phase is 5% by weight or less. If the ratio of the α-(Fe, Co, Ti) phase is high, even if the subsequent heat treatment in a vacuum atmosphere or an inert gas atmosphere is performed, there is a possibility that the main phase will not be sufficiently regenerated by recombination reaction. be. By absorbing and desorbing hydrogen, the rare earth-rich phase (subphase) expands and contracts in volume, and cracks occur between the main phase grains and between the main phase grains and the subphase. As an example, when the KHg ₂ type phase reacts with hydrogen, it changes to a phase having peaks at lattice spacing values of 3.03 Å or more and 3.10 Å or less in the powder X-ray diffraction pattern. As a result, the volume of the subphase increases and cracks occur between the grains of the main phase. During the pulverization process using jet mills, stamp mills, ball mills, etc., cracks that occur between the main phase grains or between the main phase grains and the subphase increase the probability that the magnetic powder will crack at the cracks, resulting in single crystal units. It is possible to obtain highly orientable anisotropic magnetic powder containing a large amount of fine powder of

Examples of the present invention will be specifically described below, but the present invention is not limited to these examples.

[Examples 1 to 6]
<Process A>
Raw material metals of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed in increments, taking into consideration the evaporation of rare earth elements during melting. After sufficiently melting these raw metals in the tapping pipe of a liquid ultraquenching device (melt spinning device) to form a molten alloy, the molten metal is tapped on a Cu roll rotating at a roll peripheral speed of 15 m/s. bottom. The molten metal contacted the surface of a roll rotating at high speed, was rapidly removed from the heat, extended in the form of a ribbon, and solidified. In this way, ultraquenched ribbons having the compositions shown in Table 1 were produced. As a comparative example, a superquenched ribbon having an alloy composition of Y _0.40 Sm _0.60 (Fe _0.80 Co _0.20 ) _8.33 Ti _0.70 Cu _0.63 was also produced.

<Process B>
The ultra-rapidly quenched ribbon produced in step A was wrapped in Nb foil and heat-treated at 1050° C. for 20 minutes in an Ar stream to produce alloys of Examples 1 to 6 and Comparative Example 1.

Table 1 lists the composition of the alloy, the composition of the main phase identified by SEM-EDX analysis, the saturation magnetization (Ms) at room temperature, the magnetic anisotropy field (μ ₀ H _a ), and the Curie temperature (T _c ). indicates Saturation magnetization (M _s ) and magnetic anisotropy field (μ ₀ H _a ) were evaluated using a vibrating sample magnetometer (VSM) capable of applying a magnetic field up to 10T. Since the magnetic particles are isotropic, the saturation magnetization was identified using the square-law saturation asymptotic law, and the magnetic anisotropy field was identified using the singularity detection method. The Curie temperature (T _c ) was defined as the inflection point of the curve obtained with a thermomagnetic balance. Here, the thermomagnetic balance is a device that applies a magnetic field to the sample portion of a thermobalance (TG) with a permanent magnet, and detects the temperature change of the magnetic absorption force in the sample as the weight value of the TG.

As shown in Table 1, when the amount of Ti added to the raw material alloy decreased, the Ti content of the main phase in the finally obtained alloy (after step B) decreased, while the Cu content hardly changed. . This is because Ti does not form a solid solution in the subphase, but forms a solid solution only in the main phase.

According to this example, both the saturation magnetization and the magnetic anisotropic magnetic field at room temperature generally tend to increase as the amount of Ti added decreases. Furthermore, the Curie temperature also increased as the amount of Ti added decreased. Moreover, in the comparative example in which the amount of Ti added was 0.70, the saturation magnetization (M _s ), the magnetic anisotropic magnetic field (H _a ), and the Curie temperature (T _c ) all decreased.

As examples _of structural changes before _and after the heat treatment in _step B _{, FIGS .} 4 shows the observation results of the structure of _{a 42} Cu _0.67 rare earth magnet alloy before and after heat treatment using a polarizing microscope. From FIGS. 1A and 1B, it can be seen that the microcrystals, whose size was on the order of several hundred nm before the heat treatment, grow into coarse crystal grains of 10 μm or more by the heat treatment in step B. FIG. In addition, it can be seen that there are clearly more subphases in the 1-2 phase than in the 1-4 composition phase (the black part of the grain boundary is the 1-2 phase, and the white (orange in color) linear and wedge-shaped parts are 1-4 phases).

2A and 2B show composition analysis results by SEM-EDX. From FIG. 2B, it can be seen that both the 1-2 phase and the 1-4 composition phase of the subphases have Sm richer compositions than the Y/Sm ratio of the entire alloy. Also, the 1-4 composition phase contains more Fe and Co than the 1-2 phase. Furthermore, Ti could not be measured in the detectable range in both phases.

3A and 3B are graphs showing the temperature dependence (from room temperature to 140° C.) of volume magnetization and magnetic anisotropy field of the main phase of the alloy compositions of Examples 3 to 6, respectively. However, the ratio of the main phase to the subphase was identified by X-ray Rietveld analysis, and the subphase was treated as non-magnetic. Volume magnetization decreases with increasing temperature. As shown in FIG. 3A, Examples 3 to 6 (z' in the main phase composition is 0.48 to 0.74) are higher than Comparative Example 1 (z' in the main phase composition is 0.91). volume magnetization is obtained. Further, as shown in FIG. 3B, Examples 4 to 6 (z' in the main phase composition is 0.48 to 0.67) are better than Example 5 (z' is 0.74) and Comparative Example 3 (z ' is 0.91), the decrease in the magnetic anisotropic magnetic field accompanying the temperature rise (20° C. to 100° C., which is practically important) is suppressed. Therefore, preferable ranges are 0.48≦z′≦0.67 for the main phase composition and 0.42≦z≦0.50 for the alloy composition.

[Examples 7 to 12]
<Process A>
Raw material metals of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed in increments in consideration of evaporation of rare earth elements during melting. Each element was weighed so that the amount of the main phase and the amount of the subphase shown in Table 2 were obtained. After these raw metals were sufficiently melted in a tapping pipe of a liquid ultra-quenching device (melt spinning device), the molten metal was tapped on a Cu roll rotating at a roll peripheral speed of 15 m/s. The molten metal contacted the surface of a roll rotating at high speed, was rapidly removed from the heat, extended in the form of a ribbon, and solidified. In this way, ultraquenched ribbons having the compositions shown in Table 2 were produced.

<Process B>
The ultra-rapidly quenched ribbon produced in step A was wrapped in Nb foil and heat-treated at 1050° C. for 20 minutes in an Ar stream to obtain alloys of Examples 7-12.

Table 2 also describes the presence or absence of a different phase after heat treatment with respect to the weight ratio of the main phase and the subphase. When the amount of the liquid phase was small, the foreign phase could not be eliminated. When the amount of the liquid phase was 3 wt % or more, the different phases disappeared, and an alloy for rare earth magnets having only the main phase and the subphase could be produced. FIG. 4 shows the powder X-ray diffraction results of the ultra-quenched ribbon before and after heat treatment with a subphase ratio of 3 wt %. The diffraction peaks of the TbCu ₇ phase and α-(Fe, Co, Ti) phase, which mainly existed as heterogeneous phases before the heat treatment, disappeared after the heat treatment at 1050 ° C. for 20 minutes, and the diffraction peaks of the main phase and the sub phase disappeared. was observed.

[Example 13]
Experiments were conducted in the same manner as in Examples 1 to 12 in order to observe hydrogen absorption and desorption reactions. However, the peripheral speed of the Cu roll was set to 10 m/s, and an ultra-quenched ribbon of Sm(Fe _0.16 Co _0.12 Cu _0.72 ) ₄ having a 1-4 composition as a subphase was produced. This was subjected to heat treatment up to 700° C. in a hydrogen stream, and exhaust gas was measured. That is, hydrogen absorption reduces emissions, and hydrogen release increases emissions.

FIG. 5 shows the change in the amount of hydrogen in the exhaust gas with temperature. Figure 5 shows that the subphase absorbs hydrogen in the temperature range from 250°C to 400°C and releases hydrogen in the temperature range from 540°C to 660°C. Such hydrogen absorption and release reflects that the subphase is a rare earth-rich phase relatively rich in rare earths.

[Example 14]
<Process A>
Raw material metals of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed in increments in consideration of evaporation of rare earth elements during melting. Specifically, a super-quenched ribbon having the same composition as the super-quenched ribbon produced in Example 6 was produced under the same rapid cooling conditions.
<Step B>
The ultraquenched ribbon produced in step A was wrapped in Nb foil and heat-treated at 1050° C. for 20 minutes in an Ar stream.
<Process C>
After heat treatment was performed up to 500° C. in a hydrogen stream to sufficiently absorb hydrogen, the atmosphere was switched to a vacuum atmosphere to sufficiently release hydrogen and cooled to room temperature. It was confirmed that the obtained alloy had cracks between the main phase and the subphase and in the subphase.

The rare earth magnet alloy that has undergone steps A to C and the rare earth magnet alloy that has undergone only steps A and B were pulverized by a jet mill. The degree of orientation was evaluated by orienting the obtained fine powder in a magnetic field and measuring the saturation magnetization. As a result, it was found that the rare earth magnet alloy that has undergone steps A to C has a higher magnetic flux density and a higher degree of orientation than the rare earth magnet alloy that has undergone only steps A and B. . This is mainly because cracks are introduced into the main phase and the subphase or in the subphase, thereby increasing the probability of cracking at the ends of the main phase crystal grains.

[Examples 15 to 18]
<Process A>
The alloy was made using a strip casting machine. First, raw material metals of Y, Sm, Fe, Co, Ti, and Cu having a purity of 99.9% or more were weighed in increments in consideration of evaporation of rare earth elements during melting. These starting metals were placed in a silica crucible, and the temperature was raised to 1500° C. by high-frequency induction heating in an atmosphere of −50 kPa in Ar gauge pressure to melt the starting materials. Thereafter, the molten metal was cooled to 1450° C., temporarily stored in a tundish, and then poured onto a cooling roll made of copper rotating at a peripheral speed of 1.5 m/s for cooling. The cooled alloy was pulverized by a pulverizer installed below the cooling rolls. Thus, a strip cast alloy with a composition of Y _0.36 Sm _0.64 (Fe _0.83 Co _0.17 ) _9.76 Ti _0.53 Cu _0.53 was produced. This composition was within the scope of the present invention.
<Step B>
The strip cast alloy produced in step A was placed in a molybdenum container and heat-treated at 1100° C. for 1 hour in an Ar stream.
<Process C>
After holding the alloy obtained in step B at temperature T1 (T1 = 150°C, 350°C, 600°C) in a hydrogen flow atmosphere for 1 hour, the atmosphere is switched to Ar flow atmosphere and held for 10 minutes, and then cooled to room temperature. cooled.

After the cross sections of the alloys obtained in steps B and C were embedded with resin, they were polished and processed with a cross-section polisher, and a field of view of 60 μm×45 μm was photographed with an SEM. The results are shown in Figures 7A-D. Using an image with a field of view of 60 μm×45 μm, the number of visually identifiable cracks between main phase grains was counted.

The alloys obtained in Steps B and C were pulverized in a mortar, passed through a mesh with an opening of 75 μm, and powder X-ray diffraction measurement was performed using the powder. The α-(Fe, Co, Ti) phase ratio was obtained from the obtained diffraction pattern by Rietveld analysis. DIFFRACplus Professional TOPAS 4 (manufactured by Bruker AX Co., Ltd.) was used as Rietveld analysis software. For comparison of peak heights, analysis software EVA (manufactured by Bruker AX Co., Ltd.) was used to remove the background.

In the alloy of Example 15 without hydrogen treatment, there was only one crack between the main phase grains in the field of view of 60 μm×45 μm. Even in the alloy of Example 16 after hydrogen treatment at 150° C., there was only one crack between the main phase grains. In the alloy of Example 17 after hydrogen treatment at 350° C., cracks between main phase grains were observed at 13 locations, which was significantly increased from Examples 15 and 16. In the alloy of Example 18 after hydrogen treatment at 600° C., cracks between the main phase grains were observed at three locations, and a fine structure was observed in the outer shell of the main phase grains. The α-(Fe, Co, Ti) phase ratio of Example 18 is 34.3% by weight, which is very high compared to Examples 15 to 17, and the fine structure of the main phase grain outer shell of Example 18 is mainly It can be said that the structure contains an α-(Fe, Co, Ti) phase formed by phase decomposition. As described above, when the α-(Fe, Co, Ti) phase ratio is high, even if the subsequent heat treatment in a vacuum atmosphere or an inert gas atmosphere is performed, the main phase is sufficiently regenerated by recombination reaction. Therefore, it is desirable that the α-(Fe, Co, Ti) phase ratio is low (preferably 10% or less, more preferably 5% or less).

FIG. 6 shows, from the bottom, powder X-ray diffraction patterns of samples of Examples 15 to 18, respectively, in an embodiment of the present disclosure. As shown in FIG. 6, in Examples 15 and 16 (in FIG. 6, no hydrogen treatment is Example 15, T1=150 ° C. is Example 16), the lattice spacing value is in the range of 2.50 Å or more and 2.57 Å or less. A clear KHg type ₂ (121) peak was observed in the . On the other hand, in Example 17 (T1 = 350 ° C. in FIG. 6), no clear KHg ₂ type (121) peak was observed, and instead the lattice spacing value was 3.03 Å or more and 3.10 Å or less. As a result, the peak intensity of the unidentified phase existing in the range of .sup.2 is higher (the position of .box-solid. in FIG. 6). Also, a BSE image of a field of view of 240 μm×180 μm was taken. Then, using the image analysis software Scandium, the area ratio of the subphase in the image with a field of view of 240 μm × 180 μm was calculated. 17 was 6.4.

From these results, in order to obtain highly orientable anisotropic magnetic powder containing a large amount of single crystal unit fine powder by increasing the probability of magnetic powder cracking at the crack portion, the lattice spacing value in the powder X-ray diffraction pattern should be 3. The maximum value of the X-ray intensity in the range of .03 Å or more and 3.10 Å or less is greater than the maximum value of the X-ray intensity in the range of the lattice spacing value of 2.50 Å or more and 2.57 Å or less, and It is preferable that the α-(Fe, Co, Ti) phase is 10% by weight or less (Example 17).

Since the rare earth magnet alloy of the present disclosure contains a main phase with improved magnetic properties and thermal stability and a rare earth rich sub phase, it can be suitably used to produce anisotropic sintered magnetic powder. Anisotropic sintered magnetic powder can be suitably used for producing sintered magnets. Sintered magnets are used in various motors and actuators, and have various industrial uses.

Claims (10)

An alloy for a rare earth magnet having a hard magnetic main phase having a ThMn ₁₂ type crystal structure and at least one kind of main phase not having a ThMn ₁₂ type crystal structure and a rare earth-rich secondary phase having a higher rare earth composition ratio than said main phase The composition of the entire alloy is represented by the following compositional formula (2),
R1 _1-x R2 _x (Fe _1-y Co _y ) _wz Ti _z Cu _α (2)
w, z, and α are each 8≤w≤13,
0.42≦z<0.65, and 0.46≦α≦0.70,
satisfies the
R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1,
R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is 50 mol% or more of the total of R2,
x and y are each
0.5≦x≦1.0,
0≤y≤0.4,
An alloy for rare earth magnets that satisfies The composition of the main phase is represented by the following compositional formula (3),
R1 _1-x' R2 _x' (Fe _1-y' Co _y' ) _12-z'-α' Ti _z' Cu _α' (3)
x', y', z', and α' are each
0.5≦x′≦1.0,
0≦y′≦0.4,
0.48≦z′<0.82, and
2. The rare earth magnet alloy according to claim 1 , which satisfies 0.24≤α'≤0.37. z', 0.48≤z'<0.74
The alloy for rare earth magnets according to claim 1 or 2, which satisfies 4. The alloy for a rare earth magnet according to claim 2, wherein 50 mol % or more of said entire subphase is a crystal phase containing Cu. the crystal structure of said subphase comprises a KHg _2- type crystal structure phase;
5. The rare earth magnet alloy according to claim 4, wherein the KHg ₂ -type crystal structure phase accounts for 50% or more of the subphase by volume. 6. The claim 4 or 5, wherein the subphase comprises R atoms and the R atoms present in the subphase have a molar ratio of [R2]/([R1]+[R2]) higher than the composition of the entire alloy. alloy for rare earth magnets. 7. The rare earth magnet alloy according to any one of claims 4 to 6, wherein the weight ratio of said subphase in said whole rare earth magnet alloy is 3 wt% or more and 10 wt% or less. 8. The rare earth magnet alloy according to any one of claims 1 to 7, wherein cracks are generated between said main phase and said subphase or in said subphase. The maximum value of X-ray intensity in the range of lattice spacing values of 3.03 Å or more and 3.10 Å or less in the powder X-ray diffraction pattern is 2.50 Å or more and 2.57 Å or less. 9. The rare earth magnet alloy according to any one of claims 1 to 8, wherein the X-ray intensity is higher than the maximum value and the α-(Fe, Co, Ti) phase is 10% by weight or less. 10. The rare earth magnet alloy according to claim 9, wherein the α-(Fe, Co, Ti) phase is 5% by weight or less.

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